Effect of atomic order on the martensitic transformation of Ni Fe Ga alloys

Transcription

1 Scripta Materialia 54 (2006) Effect of atomic order on the martensitic transformation of Ni Fe Ga alloys R. Santamarta a, *, E. Cesari a, J. Font b, J. Muntasell b, J. Pons a, J. Dutkiewicz c a Departament de Física, Universitat de les Illes Balears, Ctra. de Valldemossa, km 7.5, E07122 Palma de Mallorca, Illes Balears, Spain b Departament de Física i Enginyeria Nuclear, Universitat Politècnica de Catalunya, Avda. Diagonal, 647, E08028 Barcelona, Spain c Institute of Metallurgy and Materials Science, Ul. Reymonta, 25, Kraków, Poland Received 10 February 2006; received in revised form 7 March 2006; accepted 10 March 2006 Available online 3 April 2006 Abstract The effect of atomic order on martensitic transformation (MT) temperatures has been studied by differential scanning calorimetry in three polycrystalline alloys with compositions Ni 53.5+x Fe 19.5 x Ga 27.0 (x = 0, 0.5 and 1.5). Samples water quenched from different temperatures ( K) exhibit higher MT temperatures than ones slow cooled from the same temperature. This effect has been ascribed to a decrease of the L2 1 degree of order of the austenitic phase, which promotes an increase in the MT temperatures in these Ni Fe Ga alloys. The differences in ordering with cooling rate have been qualitative confirmed by electron diffraction patterns. Ó 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Shape memory alloys (SMA); Ni Fe Ga alloys; Martensitic phase transformation; Ordering 1. Introduction The ability of near-stoichiometric Ni 2 MnGa to produce up to 10% recoverable magnetic field induced strains (MFIS) [1,2], which are one order of magnitude higher than in conventional magnetostrictive materials, has attracted interest in the research and development of ferromagnetic shape memory alloys (FSMA). These large MFIS result from the high magnetocrystalline anisotropy energy of the martensite and mobility of the interfaces between martensitic variants in Ni Mn Ga alloys, in which the magneto-stress due to the applied field acts in an equivalent way to mechanical stress producing the de-twinning of martensite [3]. However, the poor ductility of Ni Mn Ga alloys could restrict the range of expected applications, motivating the development of new FSMA with better mechanical properties. * Corresponding author. Tel.: ; fax: address: ruben.santamarta@uib.es (R. Santamarta). Among the new alternatives to Ni 2 MnGa, near-stoichiometric Ni 2 FeGa alloys seem to be promising. This system shows a higher magnetization of the martensite than that of the parent phase, the former being 31.8 Am 2 /kg and 29.5 Am 2 /kg for the Ni 56 Fe 19 Ga 25 at 293 K and Ni 54 Fe 19 Ga 27 at 263 K, respectively [4,5]. The high anisotropy constant for the first alloy ( J/m 3 ) [5] is also comparable to that of Ni Mn Ga [3] and, as for the latter system, Ni Fe Ga can also exhibit ferromagnetic martensites with layered structures, which are those showing large MFIS in Ni Mn Ga alloys. The well known five-layered (10M) and seven-layered (14M) martensites [6,7] have already been observed in Ni Fe Ga [4,8] as well as the recently discovered six-layered martensite (6M) [9]. Finally, the possibility of forming small amounts of a ductile second phase (c phase) through thermal treatments is thought to improve the high brittleness of Ni Mn Ga, as occurs in Ni Al and Ni Co Al shape memory alloys [10,11]. Suitable compositions in Ni Fe Ga alloys can be selected in order to obtain, on cooling, a martensitic transformation (MT) from the L2 1 -Heusler-structure to ferromagnetic martensites around room temperature (RT) [8], which is /$ - see front matter Ó 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi: /j.scriptamat

2 1986 R. Santamarta et al. / Scripta Materialia 54 (2006) an important condition for most of the possible applications. In addition, this system can show very good thermal stability under ageing at moderate temperatures (between 520 and 770 K) [12]. In the latter work, a slight increase of the MT temperatures during the first steps of ageing was also observed, which was thought to be a consequence of changes in the degree of order. In the present work, the effect of order on the transformation temperatures of three polycrystalline Ni Fe Ga alloys was systematically studied by differential scanning calorimetry (DSC) and transmission electron microscopy (TEM) in samples either water quenched to RT or slowly cooled from temperatures in the range K. The aim of this study was to enlarge the current knowledge of the sensitivity of MT temperatures to thermal treatments in Ni Fe Ga alloys and, in particular, to establish the variation of the MT range with the degree of order. 2. Experimental procedure Three polycrystalline Ni 53.5+x Fe 19.5 x Ga 27.0 alloys with x = 0, 0.5 and 1.5, labelled as A, B and C, respectively, were cast by induction melting in argon. After solution treatment, consisting of 2 h annealing at 1270 K in an Ar atmosphere followed by slow cooling, the alloys show martensitic transformation temperatures (measured by DSC) and Curie points (T C, determined by a superconducting quantum interference device magnetometer) as detailed in Table 1. Spark cut samples from the solution-treated ingots were either water quenched or slowly cooled to RT from different temperatures, in steps of 100 K in the range K, as follows. For each temperature, a process which included three successive thermal treatments (TT) was applied to the same sample. The first one (TT1) consisted of water quenching to RT from the selected temperature, at which the sample had been kept for 20 min. The second treatment (TT2) consisted of air cooling after 20 min at the same temperature used in TT1. Finally, a third treatment identical to TT1 was applied to the sample; this thermal treatment will be labelled as TT3 to emphasise that was performed after TT1 and TT2. In this way, the occurrence of permanent changes eventually produced by quenching or slow cooling treatments could be checked. The evolution of the martensitic transformation characteristics was monitored after each step described above by differential scanning calorimetry (DSC, Setaram 92 and Perkin Elmer DSC-7). Thin foils for transmission electron microscopy (TEM) were prepared by double-jet electropolishing using a solution of 20% perchloric acid and 80% ethanol at 16 V (0.22 A) and RT. TEM observations were performed in a Hitachi H600 operating at 100 kv equipped with a (±60, ±45 ) double tilt holder. 3. Results and discussion 3.1. Calorimetric results Fig. 1(a) and (b) show the evolution of the DSC peak temperature during the first reverse transformation for alloys A and C, respectively, as a function of the TT temperature. As can be observed, the peak temperatures of the reverse transformation increased continuously as the TT1 temperature rose from 470 up to 970 K, the point at which the shift of the MT temperatures with respect to the slow cooled samples (TT2) was maximum (60 K). After a TT1 at 1070 K the peak temperature dropped at approximately the level of the one obtained for TT1 at 670 K. On the other hand, MT temperatures after TT2 did not exhibit any significant shift as the temperature of the treatment increased, and the values of the entire series were within a few degrees, very close to those of the MT after the initial annealing treatment. Finally, TT3 brought about similar values to the ones obtained by TT1, which indicates that no permanent effects were induced after this series of treatments. It is worth mentioning that the behaviour of the forward transformations is analogous to the ones exhibited in Fig. 1(a) and (b), as can be seen in Fig. 1(c), where the evolution of the peak temperature during the first forward transformation in a sample of alloy A is shown. The evolution of the peak temperatures, for both the forward and reverse transformation in samples of alloy B after the TT process, was fully equivalent to that observed for alloys A and C. It is important to notice that the martensite finish temperature, M f, for the as-annealed alloy C was slightly above RT; therefore, all the treatments in this alloy ended in the martensite phase. Furthermore, even though the alloys A and B transform below RT after the initial annealing, they undergo the MT above RT after some TT1 and TT3 treatments (see Fig. 1). Therefore, taking into account the recent confirmation that the thermal martensite stabilization effect occurs in FSMA as in Ni Mn Ga [13], martensite stabilization has to be considered as a possible mechanism for the observed shift in the first reverse transformation in quenched samples, as direct quenching to martensite is a common method to induce stabilization in Table 1 Composition, martensitic transformation temperatures and Curie points (T C ) of the studied alloys Alloy Composition (at.%) M s (K) M f (K) A s (K) A f (K) T C (K) A Ni 53.5 Fe 19.5 Ga B Ni 54.0 Fe 19.0 Ga C Ni 55.0 Fe 18.0 Ga M s : martensite start temperature; M f : martensite finish temperature; A s : austenite start temperature; A f : austenite finish temperature; T C : Curie temperature.

3 R. Santamarta et al. / Scripta Materialia 54 (2006) Fig. 2. DSC peak associated with the L2 1! B2 order disorder transition for alloy C. Fig. 1. Evolution of the peak temperature as a function of the TT temperature during the first reverse transformation for an A sample (a), a C sample (b) and during the first forward transformation for an A sample (c). conventional shape memory alloys. However, it is known that martensite stabilization only shifts the first reverse transformation, their effects being completely removed once the material retransforms to the parent phase, whereas in the studied Ni Fe Ga alloys the shift in the transformation temperatures is permanent and is also present in the direct transformation. Thus, the shift in the reverse (and direct) transformation range after quenching from different temperatures shown in Fig. 1, even in the case when the quench is carried out directly to martensite, cannot be attributed to the stabilization of this phase. The increase in the MT temperatures after TT1 and TT3 treatments has to be ascribed to a change in the degree of order of the austenite with regard to that obtained after TT2, as no other permanent changes have been observed after the series of thermal treatments. In order to better understand the different temperature increases as the quenching temperature varies, the L2 1! B2 ordering transition has been followed by DSC, and is shown in Fig. 2 for alloy C. The transition shows a peak close to 930 K and a very broad tail at lower temperatures, which indicates that the equilibrium L2 1 degree of order progressively decreases at increasing temperatures. Completely equivalent behaviour has been found for alloys A and B. Thus, a higher disorder is promoted during the holding time of the treatments performed at increasing temperatures, which is retained at RT after water quenching. Slow cooling gives sufficient time for re-ordering and reaching a degree of order, at RT, closer to the maximum allowed by the chemical composition of the alloy. The lower L2 1 degree of order retained by increasing the quenching temperature from 470 K to 970 K causes an increasing shift in the MT temperatures, as compared to slow cooled or as-annealed samples. The lower shift of the MT after quenching from 1040 K, that is from above the beginning of the second neighbour ordering transition (B2! L2 1 ), can be easily understood in terms of the higher vacancy concentration existing at these temperatures by quenching from higher temperatures. This higher vacancy concentration promotes vacancy assisted diffusion, leading to a higher degree of order than when quenching from lower temperatures. Therefore, the shift of the MT temperatures of about 60 K obtained by quenching from 970 K can be considered very close to its maximum value, according to the DSC pattern (peak temperature) of the B2! L2 1 transition. A maximum shift of 60 K was also observed in an Ni 53 Fe 20 Ga 27 alloy after annealing at various temperatures for 1 hour followed by ice water quenching [14], which can be now understood as a result of the disorder retained by the quenching process. The slight decrease of the MT peak temperatures in slowly cooled samples as the initial temperature for the TT2 increases from above 670 K (Fig. 1), again reflects the role of the increasing vacancy concentrations with temperature and its assistance in slightly improving the final L2 1 degree of order. Consequently, according to the results discussed above, the L2 1 degree of order can be modified by quenching from a broad temperature range, having a significant effect on the MT temperatures. The increase of L2 1 order of the austenitic phase promotes a decrease in the MT

4 1988 R. Santamarta et al. / Scripta Materialia 54 (2006) temperatures in these Ni Fe Ga alloys, contrary to what has been observed, for example, in Cu Zn Al alloys [15]. An additional effect, consisting of small shifts in the transformation range, concomitant with the heating runs performed in the DSC to reach temperatures above the austenite finish temperature, A f, is observed for both direct and reverse transformations. These shifts, which are due to partial recovery of the quench effects (that is, an increase in the L2 1 degree of order) are more important in alloy C, as it transforms at higher temperatures than alloys A and B. For example, the second heating in a C sample quenched from 970 K is shifted to lower temperatures by about 20 K with respect to the first heating up to 420 K; subsequent direct and reverse MT ranges remain stable, provided the maximum temperature of the DSC runs is not increased, but at temperatures 40 K higher than in slowly cooled samples. the second-neighbour ordered phase can be formed in selected regions, resulting in the appearance of the corresponding superlattice spots in EDP. In order to confirm any change in the degree of order of Ni Fe Ga samples after the two quenching rates described in the previous section, EDP were compared for samples either water quenched at RT or slow cooled from 870 K. For this experiment, alloy A was chosen from among the three Ni Fe Ga alloys studied to ensure TEM samples to be in the parent phase during observations at RT; the temperature of the TT was selected in order to obtain a large MT temperature shift in quenched samples (i.e. enhanced changes in the degree of order). It is remarkable that all the experimental EDP taken from both groups of samples exhibit L2 1 superlattice spots (Fig. 4(a)). The relative intensity of the diffracted spots has been systematically studied 3.2. TEM observations Besides the above calorimetric experiments, the change in the degree of order between water quenched and air cooled samples has also been confirmed by means of TEM observations. The existence of any ordered structure in TEM is normally detected by the presence of superlattice spots on the electron diffraction patterns (EDP) along some specific directions. For instance, in the case of Heusler alloys, the first-neighbour ordered phase (B2) can be easily distinguished from the second-neighbour ordered phase (L2 1 ) by EDP along the [110] zone axis of the parent phase. Fig. 3 shows simulated EDP along this zone axis for both perfect ordered phases in stoichiometric Ni 2 FeGa, although the intensity of the originally weak L2 1 superlattice spots has been increased for reasons of clarity. As can be observed, the existence of an extra superlattice spot at 1=2h1 11i B2 (and the equivalent ones obtained by translation symmetry) reveals the presence of L2 1 order in the parent phase of these alloys. However, it has to be noted that the perfect L2 1 structure cannot be achieved in any of the studied alloys, as the chemical composition is relatively far from the stoichiometric Ni 2 FeGa. In spite of this, Fig. 4. (a) Experimental EDP along the [110] zone axis indicating two of the four scan line profiles applied to each EDP. (b) Intensity spectrum obtained in the SLP2 from (a) in which the corresponding B2 + L2 1 peak (left) and L2 1 peak (right) are shown. Fig. 3. Simulated EDP along the [110] zone axis for B2 (a) and L2 1 (b) perfectly ordered phases.

5 R. Santamarta et al. / Scripta Materialia 54 (2006) Fig. 5. Average of all the experimental EDP taken for quenched samples (grey) and slow cooled specimens (black) after normalizing heights and rescaling. using a large number (70) of EDP taken from either water quenched or slow cooled samples, in order to increase the reliability of the results. In addition, the EDP were taken with a relatively large selected area aperture in order to obtain a good averaged region (as independent of the size of the ordered domains as possible), recorded into photographic plates and then digitalized by a scanner. Afterwards, a scan line profile (SLP) was obtained for each of the four equivalent h1 11i B2 directions connecting the 1 11-type superlattice spots (which only arise from the L2 1 structure) with the 002-type (with contributions from both L2 1 and B2 structures), as indicated in Fig. 4. All the SLP were obtained with lines of 1 pixel thickness (thicker lines were drawn in Fig. 4 for clarity reasons), crossing the maximum intensity of both peaks. Finally, all the profiles were averaged and normalized to a common height of the 002 B2+L21 peak separately for quenched samples and for slow cooled samples. The final result showed that the intensity (i.e. height) of the 1 11 L21 -type superlattice spots was about 88% the intensity of the 002 B2+L21 - type spots in the case of slowly cooled samples, whereas this ratio was about 63% for the water quenched specimens (Fig. 5). In this figure the B2 + L2 1 -type and L2 1 -type peaks have also been shifted along the x axis for clarity reasons. Therefore, the TEM experiments clearly confirm, in a qualitative way, that quenched samples show a lower L2 1 degree of order than slow cooled samples. 4. Conclusions Thermal treatments consisting of water quenching or air cooling from different temperatures in the range K have been performed for three polycrystalline Ni Fe Ga alloys. In these alloys, the lower L2 1 degree of order retained by water quenching as compared to slowly cooled samples, promotes higher MT temperatures. Increasing disorder can be retained by quenching from increasing temperatures up to about K, where a maximum shift close to 60 K in the MT range can be obtained. Quenching from higher temperatures does not increase the retained disorder as the higher vacancy concentrations assist in the ordering reaction. The differences in the degrees of order have been confirmed by statistical analysis of intensities of B2 and L2 1 superlattice spots in electron diffraction patterns. Acknowledgements Partial financial support from the Spanish DGI, research project MAT , is acknowledged. Dr. J.I. Pérez- Landazábal and Dr. V. Recarte are also acknowledged for the T C measurements. References [1] Ullakko K, Huang JK, Kanter C, Kokorin VV, O Handley RC. Appl Phys Lett 1996;69:1966. [2] Müllner P, Chernenko VA, Kostorz G. J Appl Phys 2004;95:1531. [3] Chernenko VA, L vov VA, Cesari E. J Magn Magn Mater 1999; :859. [4] Oikawa K, Ota T, Sutou Y, Ohmori T, Kainuma R, Ishida K. Mater Trans 2002;43:2360. [5] Li Y, Jiang C, Liang T, Ma Y, Xu H. Scripta Mater 2003;48:1255. [6] Pons J, Chernenko VA, Santamarta R, Cesari E. Acta Mater 2000;48:3027. [7] Pons J, Santamarta R, Chernenko VA, Cesari E. J Appl Phys 2005;97: [8] Oikawa K, Ota T, Ohmori T, Tanaka Y, Morito H, Fujita A, et al. Appl Phys Lett 2002;81:5201. [9] Pons J, Santamarta R, Chernenko VA, Cesari E. Mater Sci Eng A, in press. [10] Ishida K, Kainuma R, Ueno N, Nishizawa T. Metall Trans A 1991;22:441. [11] Oikawa K, Wulff L, Iijima T, Gejima F, Ohmori T, Fujita A, et al. Appl Phys Lett 2001;79:3290. [12] Santamarta R, Font J, Muntasell J, Masdeu F, Pons J, Cesari E, et al. Scripta Mater 2006;54:1105. [13] Seguí C, Cesari E, Font J, Muntasell J, Chernenko VA. Scripta Mater 2005;53:315. [14] Omori T, Kamiya N, Sutou Y, Oikawa K, Kaimuna R, Ishida K. Mater Sci Eng A 2004;378:403. [15] Rapacioli R, Ahlers M. Acta Metall 1979;27:777.